Metal composites and methods for forming same

ABSTRACT

A metal composite comprising a spinodal structure having at least one ductile phase and method of making same is disclosed. The metal composite is formed by forming an alloy comprising a positive heat of mixing in the liquid state; purifying the alloy; and forming a network structure of the alloy comprising at least one ductile sub-network.

RELATED APPLICATIONS

This application claims the benefit under 35 U.S.C. §120 of U.S.application Ser. No. 11/088,106, titled “METAL COMPOSITES AND METHODSFOR FORMING SAME,” filed on Mar. 23, 2005, which is herein incorporatedby reference in its entirety.

FIELD OF INVENTION

This invention relates generally to a metal composite and methods forforming same and, more particularly, to a metal composite having anetwork structure with at least one ductile sub-network.

BACKGROUND

Nanostructured materials are defined as materials having a grain sizediameter of between and including 1 nm and 1,000 nm. The presence ofnanostrutures in a material improves mechanical properties when comparedto the same material formed without nano structures.

Nanostructured materials have typically been synthesized by powdersintering or thermal annealing of glassy metals, and fluxing. In powdersintering, nanometer size powders are compacted together, and may beannealed, to produce nanostructured materials. Nanostructured materialsproduced by powder sintering are typically disk shaped, having adiameter of about 1 cm, and a thickness of 1 mm to 2 mm. The powderedsintered materials are typically brittle and exhibit voids andnon-uniform grain growth.

Alternatively, thermal annealing of an amorphous metal producesnanocrystals in an amorphous matrix. A metallic melt is quenched into aglass or amorphous metal, which is then annealed at a temperature nearits glass transition temperature, resulting in uniformly distributednanocrystals in an amorphous matrix. When the amorphous matrix alsocrystallizes, a crystalline nanostructure may be produced.

A fluxing technique has more recently been used to preparenanostructured materials. In this method, liquid-state spinodaldecomposition occurs resulting in brittle solid spinodals. A typicalPd—Si nanostructure is shown in FIG. 1. As seen in FIG. 1, discrete Pdprecipitates (shown in white) are located inside solid spinodals.However, the discrete Pd precipitates are at a negligible volumefraction so that the solid spinodals remain brittle.

As such, the major constituent phases of conventional nanostructures arebrittle, and have a broad grain size distribution. Conventionalnanostructures also have low strength and low impact fracture energy.Moreover, conventional to amorphous metals typically have a smalloverall dimension, such as ribbon or foil forms having a thickness of30-50 microns that renders them inappropriate for commercialapplications.

SUMMARY OF INVENTION

In one embodiment, the present invention is directed to a methodcomprising forming an alloy with the constituents, purifying the alloy,and forming a network structure of the alloy comprising at least oneductile sub-network structure.

Another embodiment is directed to a metal composite comprising a ductilespinodal structure. Another embodiment is directed to a metal articlecomprising the metal composite. In another embodiment, the metal articleis a nanostructure composite.

Another embodiment is directed to a method of forming a metal compositecomprising forming an alloy, purifying the alloy, forming one or morespinodals, and heating the one or more spinodals causing at least one ofthe one or more spinodals to transform into one or more ductile phases.

Other advantages, novel features, and objects of the invention willbecome apparent from the following detailed description of non-limitingembodiments of the invention when considered in conjunction with theaccompanying figures. In cases where the present specification and adocument incorporated by reference include conflicting disclosure, thepresent specification shall control.

BRIEF DESCRIPTION OF DRAWINGS

In the figures:

FIG. 1 is a TEM micrograph of a conventional Pd—Si nanostructure at amagnification of 75,000×.

FIG. 2 is a SEM micrograph showing the microstructure of a conventionalsteel ball at a magnification of 9,500×.

FIG. 3A is a SEM micrograph showing a portion of a two phase spinodalmicrostructure of Fe₈₀C₁₅Si₅ at a magnification of 35,000×.

FIG. 3B is a SEM micrograph showing another portion of the two phasespinodal microstructure of FIG. 3A at a magnification of 6,000×.

FIG. 3C is a SEM micrograph showing another portion of the two phasespinodal microstructure of FIG. 3A at a magnification of 34,000×.

FIG. 3D is a SEM micrograph showing another fracture surface of the twophase spinodal microstructure of FIG. 3A at a magnification of 8,500×.

FIG. 3E is a SEM micrograph showing a fracture surface of the two phasespinodal microstructure of FIG. 3A at a magnification of 9,400×.

FIG. 4A is a SEM micrograph showing a portion of a spinodal structure ofFe_(40.5)CO_(40.5)C₁₄Si₅ at magnification of 34,000×.

FIG. 4B is a SEM micrograph showing another portion of the spinodalstructure of FIG. 4A at a magnification of 5,400×.

FIG. 4C is a SEM micrograph showing a fracture surface of the spinodalstructure of FIG. 4A at a magnification of 13,500×.

FIG. 4D is a SEM micrograph showing another fracture surface of thespinodal structure of FIG. 4A at a magnification of 4,100×.

FIG. 5A is a SEM micrograph showing a portion of a spinodal structure ofCo₇₅Si₁₅B₁₀ at a magnification of 34,000×.

FIG. 5B is a SEM micrograph showing another portion of the spinodalstructure of FIG. 5A at a magnification of 2,000×.

FIG. 5C is a SEM micrograph showing another portion of the spinodalstructure of FIG. 5A at a magnification of 5,400×.

FIG. 5D is a SEM micrograph showing a fracture surface of the spinodalstructure of FIG. 5A at a magnification of 1,400×.

FIG. 6A is a SEM micrograph showing a portion of a spinodal structure ofFe₈₂C₁₈ at a magnification of 4,100×.

FIG. 6B is a SEM micrograph of a fracture surface of the spinodalstructure of FIG. 6A at a magnification of 9,500×.

DETAILED DESCRIPTION

This invention is not limited in its application to the details ofconstruction and the arrangement of components set forth in thefollowing description or illustrated in the figures. The invention iscapable of other embodiments and of being practiced or of being carriedout in various ways. Also, the phraseology and terminology used hereinis for the purpose of description and should not be regarded aslimiting. The use of “including,” “comprising,” or “having,”“containing,” “involving,” and variations thereof herein, is meant toencompass the items listed thereafter and equivalents thereof as well asadditional items.

The present invention provides a metal composite comprising a network ofat least two spinodals, or sub-networks, and methods for forming same.The terms “spinodal” and “sub-network” are used interchangeably herein,to define a solid morphology of isolated clusters and/or interconnectedregions following liquid state spinodal decomposition and subsequentsolidification. For example, spinodal decomposition and subsequentsolidification of a binary alloy results in two sub-networks. Similarly,spinodal decomposition of a tertiary alloy results in two or threesub-networks. Sub-networks may include discrete precipitates, but not inan appreciable amount.

In one embodiment, at least one spinodal or sub-network of the metalcomposite is a ductile phase. As used herein the phrase “ductile phase”is defined as a malleable phase. The ductile spinodal may be isolatedclusters, partially interconnected, substantially interconnected, andcombinations thereof. That is to say that the degree of interconnectionof the ductile spinodal may vary throughout the metal composite. In oneembodiment, the metal composite comprises a ductile spinodal and abrittle spinodal. In another embodiment, the metal composite tocomprises a first ductile spinodal and a second ductile spinodal. In oneembodiment of the invention, relative volume fractions of the ductileand brittle phases may be changed to vary any desired property, such asbut not limited to, hardness, fatigue strength, compressive yieldstrength, wear resistance, and maximum operating temperature.

The metal composite may be formed from any alloy for which a predominantconstituent has a positive heat of mixing with at least one otherconstituent in the liquid state. As used herein, the phrase “heat ofmixing” is defined as an enthalpy change when 1 mol of a mixture isformed from its pure components at a temperature, T. As used herein, theterm “predominant” is defined as an intended primary constituent of thealloy. The liquid state may be either stable or metastable. Examples ofalloys having constituents with a positive heat of mixing of include:monotectic, eutectic, and peritectic alloys. Other examples of alloyshaving constituents with a positive heat of mixing are noted in the text“Thermodynamics of Solids,” by Richard A. Swalin, Published by JohnWiley & Sons, Inc. (1962), which is incorporated herein by reference forall purposed. In one embodiment, the alloy may comprise a metal and ametalloid.

In one embodiment, the overall heat of mixing of all constituents in thealloy may be positive. In another embodiment, the overall heat of mixingin the liquid state of all constituents in the alloy may be negative.For example, an alloy comprised of first, second, and third componentsmay have an overall negative heat of mixing in the liquid state,although the first constituent (the major constituent) has a positiveheat of mixing in the liquid state with the second component, and alsohas a positive heat of mixing in the liquid state with the thirdcomponent. In this example, an overall negative heat of mixing may arisefrom a very negative heat of mixing in the liquid state for the secondand third constituents, which overwhelms the positive heats of mixing inthe liquid state for the first major constituent with the second andthird constituents.

Any metal may be used to obtain desired composite properties as long asthe selected metal has a positive heat of mixing when combined with atleast one other constituent of the alloy in the liquid state. Forexample, the metal may be a Group VIII metal, such as Fe, Co, Ni, Ru,Rh, Pd, Os, Ir, Pt, and combinations thereof. In one embodiment, themetal may be selected from the group consisting of Fe, Co, Ni, Cu, Pd,Pt, Mn, Al, T, Zr, Cr, W, and combinations thereof. In one embodiment,the metal is Co. In another embodiment, the metal is Fe. In yet anotherembodiment, the metal is Ni.

Any metalloid may be used to obtain desired composite properties, aslong as the selected metalloid has a positive heat of mixing whencombined with other predominant constituents in the liquid state. Forexample, the metalloid may be any of B, C, Si, As, Sb, Te, Po andcombinations thereof. In one embodiment, the metalloid may be any of B,C, Si, and combinations thereof. In another embodiment, the metalloid isC.

In another embodiment, the alloy comprises Fe and at least one of Si andC. In this embodiment, Fe may range between and including about 70atomic % and about 92 atomic %, Si may range between and including about0 atomic % and about 20 atomic %, and C may range between and includingabout 0 atomic % and about 30 atomic %. The ranges used representminimums and maximums of individual components, though it is understoodthat individual components are combined in such a way to result in anatomic percent of the alloy of 100%. It is believed that these ranges ofFe, Si, and C will result in a ductile spinodal network when processedaccording to the methods disclosed herein. Although compositions outsidethese ranges may produce the desired ductile nanostructure, it isbelieved these ranges may be more effective under processing conditions.

In one embodiment, the Fe—C—Si composition may range between andincluding Fe₇₆C₂₄Si₀; Fe₈₁C₁₉Si₀; Fe_(85.5)C₀Si_(14.5);Fe_(88.5)C₀Si_(11.5) and all compositions in between. It has been foundthat this range, and all points therein, result in a nanostructurecomprising ductile spinodals upon air cooling, and substantialundercooling.

In another embodiment, the Fe—C—Si composition may range between andincluding Fe₈₁C₁₉Si₀; Fe₈₄C₁₆Si₀; Fe₉₀C₀Si₁₀; Fe_(88.5)C₀Si_(11.5) andall points in between.

In yet another embodiment, the Fe—C—Si composition of another embodimentranges between and including Fe₇₃C₂₇Si₀; Fe₇₆C₂₄Si₀; Fe₈₄C₀Si₁₆;Fe_(85.5)C₀Si_(14.5) and all points therein. It has been found thatwithin these ranges, and all points there between, portions of thecomposite formed from these compositions form nanostructures comprisingductile spinodals upon air cooling. However, increased cooling rates maybe used so that the entire sample of these compositions may form thedesired nanostructures.

In another embodiment, the Fe—C—Si composition may range between andincluding Fe₈₄C₁₆Si₀; Fe₈₇C₁₃Si₀; Fe₉₀C₀Si₁₀; Fe₉₂C₀Si₈ and all pointsin between. In another embodiment, the Fe—C—Si composition of anotherembodiment ranges between and including Fe₇₀C₃₀Si₀; Fe₇₃C₂₇Si₀;Fe₈₂C₀Si₁₈; Fe₈₄C₀Si₁₆ and all points therein. It is believed that withthese ranges, the desired nanostructures may form in a drop tower filledwith a gaseous medium.

In another embodiment, B may be added to the Fe—C—Si alloy. For example,B may be added in a range of between and including about 0 atomic % andabout 5 atomic % without significantly impacting the formation of thespinodal structures.

In yet another embodiment, the alloy may comprise any non-metal selectedfrom the group Ge, P, S and combinations thereof. In one embodiment, Gemay replace, or be used in addition to Si, without significantlyaffecting the nanostructure, However, in some instances, the presence ofGe resulted in the formation of voids. In another embodiment, P may beadded to the alloy to increase the formation of spinodal structures. Forexample, between and including about 0.5 atomic % and about 4 atomic %of P may be added to the Fe—C—Si alloy so that the entire composite maycomprise a spinodal structure. It has been found that the addition of Pin the Fe—C—Si alloy reduces or eliminates the presence of eutecticstructures, thereby increasing the amount of spinodal structures.

In another embodiment, Ni may be added to the Fe—C—Si ally to increasethe volume fraction of the ductile phase. For example, about 1 atomic %to more that about 3 atomic % Ni may be added to the Fe—C—Si alloy toincrease the volume fraction of the ductile phase.

The alloy may have a glass forming ability (GFA) defined herein as aratio of a glass transition temperature (T_(g)) to a liquidustemperature (T_(l)), greater than or equal to about 0.35 so that thenanostructures formed are sufficiently large to impart the desiredphysical properties to the composite. In one embodiment, the GFA isgreater than or equal to about 0.49. For example, Fe_(82.5)B_(17.5) hasa GFA of about 0.35; Fe₈₀B₂₀ has a GFA of about 0.49; Fe₈₀C₇P₁₃ has aGFA of about 0.58; Fe₇₉Si₁₀B₁₁ has a GFA of about 0.58, and Co₇₅Si₁₅B₁₀has a GFA of about 0.56.

The alloy may be formed by heating a selected composition ofconstituents in a desired proportion. Heating may be carried out undernormal alloying conditions with conventional equipment, such as a radiofrequency induction furnace or high temperature furnace.

Optionally, the formed alloy may be divided into smaller portions of thealloy for further processing. The alloy may be placed in a first portionof a container which also has a second portion smaller than the firstportion. A vacuum may drawn on the container. The container may beheated allowing the alloy to melt forming a first molten alloy. Thefirst molten alloy may be forced into the second portion of thecontainer with a pressurized gas. The container and the first moltenalloy may be cooled forming a solidified alloy. The solidified alloy maybe removed from the second portion of the container and may beapportioned into a desired size or mass.

In one embodiment, the first portion of the container may have a crosssectional area smaller than a cross sectional area of the secondportion. Accordingly, the second portion may be sufficiently long toaccommodate the entirety of the molten alloy in the narrower crosssectional area. Upon cooling the first molten alloy and the container,the solidified alloy may be cut into various thicknesses, suitable for adesired application.

The alloy not further divided or the solidified alloy further dividedinto smaller portions may be further processed to remove impuritiesaccording to conventional methods. For example, the solidified alloy maybe heated to a temperature greater than or equal to its liquidus (T_(l))in the presence of a flux to form a second molten alloy. In oneembodiment, the solidified alloy and flux material may be heated to atemperature greater than about 1,000° C.

Any flux may be used so long as it does not react with the second moltenalloy. For example, the flux may be boron oxide, glass, calcium oxide,barium oxide, aluminum oxide, magnesium oxide, lithium oxide, andmixtures thereof. In one embodiment, the flux is glass. In anotherembodiment the flux is boron oxide. Boron oxide (B₂O₃), nominallyanhydrous, is available from Atomergic Chemetals Corporation(Farmingdale, N.Y.).

The second molten alloy may then be cooled to a temperature sufficientto form a second solidified alloy. In one embodiment, the second moltenalloy may be undercooled by cooling the second molten alloy below itsliquidus. The second molten alloy may be cooled to or below a criticaltemperature, T_(C), typically below the liquidus to allow liquid statespinodal decomposition, thereby forming liquid spinodals. Without beingbound by any particular theory, it is believed that the fluxing processpurifies the second molten alloy, allowing the second molten alloy toremain in its liquid state well below it liquidus. By allowing thesecond molten alloy to remain in its liquid state below its liquidus,the second molten alloy is a metastable liquid that undergoes spinodaldecomposition upon entering the metastable miscibility gap (criticaltemperature T_(C)), which is often substantially below the liquidus ofthe alloy. During spinodal decomposition, the molten alloy splits into anumber of metastable liquid spinodals having a liquid phase wavelengthλ. As used herein, the wavelength (λ) is defined as a lateral dimension,or diameter, of the spinodal. In one embodiment, the metastable liquidmay have a liquid phase λ of less than about 300 nm, preferably lessthan about 100 nm.

Upon cooling, the second molten alloy solidifies to form an undercooledspecimen having a spinodal or sub-network structure. The solid spinodalsmay be crystalline, amorphous, quasi-crystalline, and mixtures thereof.In one embodiment, the alloy is cooled to a ΔT of about 100° K to about500° K. As used herein the term ΔT is defined as the difference betweenthe liquidus temperature and actual temperature (T_(l)−T).

The solid phase wavelengths (λ) formed by the liquid state spinodaldecomposition of the second molten alloy typically range from microns tonanometers. The resultant composite may have a fine microstructuredefined as a material having a grain size diameter, or wavelength, ofbetween and including 1 nm and 100,000 nm. The composite may includenanostructures, wherein one physical dimension of one constituent phaseis about 1,000 nm or less. Each spinodal, or sub-network, within theentire networked structure my have solid phase wavelengths that aresimilar to or differing from one another.

In one embodiment, the spinodal or sub-network structure formed may havea solid phase λ of less than about 50 microns. In another embodiment,the spinodal or sub-network structure formed may have a solid phase λ ofabout 10 microns or less. In yet another embodiment, the spinodalstructure formed has a solid phase λ of about 300 nm or less, preferablyless than about 100 nm.

The solid phase wavelength may vary throughout a specimen. For example,the solid phase wavelength may change during crystallization. Duringcrystallization, a crystallization front moves across the molten so thatthe solid phase λ increases, effectively creating a coarser spinodalstructure. Therefore, the solid phase wavelength may be smallest at asite where crystallization is initiated, and increase as crystallizationproceeds further from the initiation site. The wavelength of the solidspinodal at the site where crystallization is initiated may be similarto the wavelength of the liquid spinodal. In one embodiment, thespinodal morphology at a distance from the site where crystallization isinitiated may be replaced (partially or entirely) by other structures,including dendrite and eutectic.

If crystallization is bypassed after the metastable liquid alloyundergoes metastable liquid spinodal decomposition, the liquid spinodalsmay solidify into amorphous spinodals upon further cooling. Thesolidification may occur homogeneously, that is, hardening does notbegin at any single location. The hardening continues on furthercooling, until all the liquid spinodals become amorphous solids. It isexpected that the solid phase λ is substantially uniform in this mode ofsolidification, so that the liquid phase λ of the liquid spinodals maybe to similar to the solid phase λ of the amorphous spinodals. Solidspinodals may also form as a mixture of amorphous and crystallinespinodals, so that λ may vary through out the solid.

The phases in the spinodals, if crystalline, may, but need not, form amicrostructure comprising a coherent grain boundary. As used herein, theterm “coherent grain boundary” is defined as a coherent interface and/ora semi-coherent interface. A coherent interface has interfacial energiesof about 10 to about 100 mJ/m², and occurs when two crystals perfectlymatch at an interface plane so that the two lattices are continuousacross the interface. A semi coherent interface occurs when theinterface has a series of edge or screw dislocations and has aninterfacial energy of about 200 mJ/m² to about 500 mJ/m².

In another embodiment of the invention, any brittle spinodals present ina specimen may be further treated to undergo phase transformation intoone or more ductile spinodals. For example, on annealing, Fe₃Sidecomposes into Fe and graphite, thereby further increasing the strengthand impact fracture of the specimen. In one embodiment, one or morebrittle spinodals may be annealed to form ductile phases. The resultantductile phases may be isolated clusters, partially or substantiallycompletely interconnected, and combinations thereof. The ductile phasesmay be interconnected with other ductile phases, and or with brittlephases.

The metal composite may be a bulk material having any shape suitable fora particular purpose. As used herein, the phrase “bulk material” isdefined as a material having a shape with a cross sectional dimensiongreater than or equal to about 1 mm in all directions. For example, thecomposite may be a sphere, cone, pyramid, square, rectangle or irregularshape. In one embodiment, the composite is a sphere. In anotherembodiment, the bulk material has at least one cross sectional dimensionof about 2.54 cm, preferably about 1 cm. In yet another embodiment, themetal composite may be a sphere having any diameter suitable for aparticular purpose. For example, the metal composite may be a spherehaving a diameter less than about 1 inch, less than or equal to about 2cm, less than or equal to about 1.0 cm, and less than or equal to about5 mm respectively. In another embodiment, the spherical metal compositemay have a diameter of about 0.1 mm. The process for forming the metalcomposite into spheres, such as ball bearings, is advantageous whencompared to conventional process of making steel balls. For example, thetypical expensive process may be replaced with a simple and inexpensivefluxing process, and the conventional heat treating process may beeliminated.

EXAMPLES

The invention may be further understood with reference to the followingexamples, which are intended to serve as illustrations only, and not aslimitations of the present invention as defined in the claims herein.

Alloys for each of the compositions listed below were prepared asfollows:

A desired composition of raw materials, selected from Fe, Co, Ni, C, Si,B, Ge, and P, were alloyed in an RF induction furnace at a minimumtemperature of about 1,000° C. The alloy was air cooled to solidify, andthen positioned in a large portion of a fused silica tube. The fusedsilica tube comprised a large portion with an inner diameter of about 10mm to 30 mm, and a long small portion having an inner diameter of about2 mm to about 8 mm. The small portion had a length of about 10 mm toabout 600 mm. The silica tube containing the alloy was evacuated by amechanical pump and placed in a furnace at a sufficient temperature andfor a sufficient time to melt the alloy. When the alloy was completelymelted, a pressurized gas was introduced into the large portion of thesilica tube forcing the molten alloy into the small portion of the tube.The tube and alloy were cooled, forming a rod shaped alloy. The rodshaped alloy was removed from the tube and cut into smaller disk shapedpieces each having a thickness in the range of between and includingabout 1 mm to about 10 mm.

Each disk was positioned with anhydrous B₂O₃ in individual fused silicatubes, having inner diameters between about 3 mm and about 15 mm, and alength of about 10 mm to about 100 mm A number of fused silica tubescontaining an alloy disk and anhydrous B₂O₃ were placed in a largerfused silica tube having a diameter of about 20 mm to about 100 mm Avacuum was drawn on the larger silica tube, which created a vacuum inthe individual tubes containing each alloy disk and anhydrous B₂O₃. Thelarge tube was then heated at sufficient temperature and for a specifiedtime period, ranging from about 15 minutes to about 8 hours tocompletely melt the alloys. The molten alloys were cooled and allowed tocrystallize at the ΔT temperatures noted below.

Example I. Fe: 80 atomic %

-   -   C: 15 atomic %    -   Si: 5 atomic %

This alloy was prepared by purifying molten Fe₈₀C₁₅Si₅ in a flux aboveits liquidus, and subsequent undercooling below its liquidus.

The Fe—C—Si system was formed into a precision ball bearing with some ofits properties listed in Table 1. Also listed in Table are comparativeresults for conventional chrome steel balls (available from FAG Bearing,Danbury, Conn.).

TABLE 1 Chrome steel balls Fe—C—Si (available from FAG Property systemBearing, Danbury, CT) Hardness (HV) 750-850 Fatigue strength >3,600**<2,600*** (Maximum compressive pressure (MPA))* Compressive yield About7,000 About 3,600 strength (MPA) Maximum operating 550° C. 150° C.temperature *Fatigue strength was determined by using a cycliccompressive force, with a minimum compressive pressure of bout 0 MPA anda maximum cyclic compressive pressure of about 3,600 MPA for 10⁷ cycles.**Samples remained in tact after completion of the test. ***Samples didnot survive the test.

It was found that the Fe—C—Si system had about the same wear resistanceand hardness as the conventional steel ball. However, balls made fromthe Fe—C—Si system exhibited a compressive yield strength almost doublethat of the conventional steel ball (7,000 MPA compared to 3,600 MPA).The balls made from the Fe—C—Si system also exhibited a greater fatiguestrength of 3,600 MPA without failure, compared to the conventionalballs which failed at 2,600 MPA. Similarly, the balls of the Fe—C—Sisystem exhibited a greater Young's modulus and greater thermal stabilitythan the conventional steel balls. In addition, the Fe—C—Si system ballsexhibited an impact fracture energy that approached that of a ceramicprecision ball bearing made of SiN.

FIG. 3A is a SEM of a two phase spinodal microstructure of theFe₈₀C₁₅Si₅ ball. The Fe₈₀C₁₅Si₅ was analyzed at various locations withinthe ball. The micrograph in FIG. 3A shows a two-phase spinodal structurehaving an interconnected microstructure. Both phases are crystalline,and have an average wavelength of about 300 nm. The randomness of themicrostructure indicates that this is the site where crystallization wasinitiated. FIG. 3B is a micrograph of a center of the specimen of FIG.3A. FIG. 3B shows alignments and differing orientation of the alignmentsof the spinodal structure near the center of the specimen. FIG. 3C is amicrograph of the specimen at an end opposite to the site of initiation.FIG. 3C shows the formation of eutectic structures at the end oppositeto the site of initiation. FIG. 3D shows the fracture behavior of thissystem, in which the fracture surface is scale-like or cloud-like. Thewhite curves in FIG. 3D are ridges, illustrating that ductile fracturehas occurred. FIG. 3E is another micrograph of the specimen of FIG. 3Ashowing the microstructure on a fracture surface of the specimen. Thetwo solid spinodals in the metal composite are Fe₃Si and body centeredcubic (BCC) Fe, or solid solution of Fe. The former spinodal is brittle,while the latter is ductile. Upon fracturing, the ductile spinodal formsridges on a fracture surface. Without being bound by any particulartheory, it is believed that the high strength and high impact fractureenergy are due to BCC Fe (or solid solution Fe).

At a very large ΔT, for example 100-500° K, the melt split into twoliquid spinodals. These two liquid spinodal were metastable and were,therefore, prone to crystallization. Crystallization started at a pointon the surface of the molten specimen (called the site of initialcrystallization). The crystallization front then spread out until theentire molten specimen became a crystal. During crystallization, heatwas released (heat of crystallization) that may have been partlyresponsible for the change in microstructures at increasing distancesfrom the site of initial crystallization. At the site of initialcrystallization, since crystal growth was fast and heat released isrelatively small, the morphology of the liquid spinodals may have beensimilar to the morphology of the solid spinodals. Moving away from thesite of initial crystallization, the spinodal morphology evolved.

In comparison, FIG. 2 is a SEM of a microstructure of a conventionalsteel ball, typically a mixture of austenite and martensite, produced byconventional methods. Typically, steel in the form of a wire coil (stillaustenite, which is soft) is manufactured. Short cylindrical pieces arecut form the wire and cold forged in a heading machine. The surface ofthe headed balls are ground by a flashing machine. The balls are thenhardened in a furnace, transforming more than half of the austentiteinto martensite, which is hard. After hardening, the balls go though twomore grinding processes (grinding and lapping) to produce a desiresurface finish. The balls are then cleaned and polished. Drawbacks tothe conventional method of preparing steel balls include the necessityof using a high quality wire coil, and ensuring the heat treatment doesnot convert all the austenite into martensite. Moreover, because not allthe austenite is converted, the conventional steel balls may not besuitable for high temperature application, such as those greater than150° C.

Although not being bound by any one particular theory, it is believedthat the microstructure including the ductile phase of the presentinvention provides a micro- or nano-structured object with uniquephysical properties when compared to similar objects made withconventional methods, such as: greater compressive yield strength,greater Young's modulus, greater fatigue resistance, and greater thermalstability, while maintaining similar wear resistance and hardness, whilealso exhibiting similar impact fracture energy similar to that ofceramic balls.

Example II. Fe: 40.5 atomic %

-   -   CO: 40.5 atomic %    -   C: 14 atomic %    -   Si: 5 atomic %

The crystallization temperature was about 800° C. The microstructures ofan as formed Fe_(40.5)CO_(40.5)C₁₄Si₅ ingot (the specimen) are shown inFIGS. 4A-4D. FIG. 4A shows the microstructure of a portion of thespecimen that crystallized first, locating the free surface of thespecimen. In FIG. 4A, there are two solid spinodals, of differentphases, which together form a random network. FIG. 4B shows themicrostructure of a portion of the specimen that crystallized last, atthe opposite end of the specimen. As seen in FIG. 4 b, themicrostructure is a mixture of spinodal morphology and eutecticstructure. FIG. 4 c shows the fracture surface of the specimen, in whichbright ridges indicate ductile fracture had take place in one of the twospinodals. The distribution of the ridges is consistent with theprediction of a spinodal mechanism. In regions dominated by eutectics,the eutectic morphology is exhibited on the fracture surface shown inFIG. 4D.

Example III. CO: 75 atomic %

-   -   Si: 15 atomic %    -   B: 10 atomic %

Crystallization occurred at about 800° C. The microstructures of an asformed CO₇₅Si₁₅ B₁₀ ingot (specimen) are shown in FIGS. 5A-5D. FIG. 5Ashows the microstructures of a portion of the specimen whichcrystallized first, located at the free surface of the ingot. In FIG.5A, there are two solid spinodals of different phases, forming anetwork. The network morphology changes as the location in the specimenmoves away from the site where crystallization was initiated. Toward thecenter of the specimen, one solid spinodal shown apparently breaks intoelongated grains, as shown in FIG. 5B. The elongated grains may be overabout 20 microns. A layer of multiple phases separates the elongatedgrains. FIG. 5C shows microstructures at a location further from thesite of initial crystallization. As shown in FIG. 5C, the layersurrounding the elongated grains of the portion of the specimen thatcrystallized last appear to be substantially homogeneous. FIG. 5D showsa fracture surface of the specimen which illustrates that the layerssurrounding the elongated grains are ductile. Without being bound to anyparticular theory, it is believed that the ductile layers provide thespecimen with impact fracture energy and strength.

Example IV. Fe: 82 atomic %

-   -   C: 18 atomic %

Crystallization temperature was about 650° C. The microstructure of anas prepared ingot of Fe₈₂C₁₈ specimen is shown in FIG. 6A. In FIG. 6A,there are many domains, each of which is occupied by a network-likestructure that is aligned. The boundaries between the domains arerelatively flat, and the alignments of the network close to theboundaries are sharp. Apparently, a form of rearrangement of the networkbranches occurred during crystallization forming the boundaries. FIG. 6Bshows the microstructure of a fracture surface of the specimen of FIG.6A. As seen in FIG. 6A, two vertical lines are found near the center ofthe micrograph, and represent the boundaries described above. Attachedto these boundaries are aligned features, similar to that of a dendrite.Further away from the boundaries, scale-like or cloud-like structuresagain dominate the microstructure. The lighted parts of FIG. 6Brepresent ridges which have undergone a ductile fracture.

Having thus described several aspects of at least one embodiment of thisinvention, it is to be appreciated various alterations, modifications,and improvements will readily occur to those skilled in the art. Suchalterations, modifications, and improvements are intended to be part ofthis disclosure, and are intended to be within the spirit and scope ofthe invention. Accordingly, the foregoing description and drawings areby way of example only.

What is claimed is:
 1. A method comprising: forming an alloy; purifyingthe alloy; and forming a network structure of the alloy comprising atleast one ductile sub-network structure.
 2. The method of claim 1wherein the alloy formed comprises a ratio of T_(g) to T_(l) greaterthan or equal to about 0.35.
 3. The method of claim 2, wherein the alloyformed comprises a ratio of T_(g) to T_(l) greater than or equal toabout 0.49.
 4. The method of claim 2, wherein the alloy formed comprisesa metal and a metalloid.
 5. The method of claim 4, wherein purifying thealloy comprises: heating the alloy to form a molten alloy; andcontacting the molten alloy with a flux material.
 6. The method of claim2, wherein forming a network structure comprises cooling the moltenalloy.
 7. The method of claim 6, wherein cooling the molten alloycomprises undercooling the molten alloy.
 8. The method of claim 7,wherein the molten alloy is cooled to a ΔT of about 100° K to about 500°K
 9. The method of claim 4, wherein the alloy formed comprises a metalselected from the group consisting of Fe, Co, Cu, Ni, Pd, Pt, Mn, Al,Ti, Zr, Cr, W, and combinations thereof.
 10. The method of claim 9,wherein the alloy formed comprises a metal selected from the groupconsisting of Fe, Co, Ni, and combinations thereof.
 11. The method ofclaim 10, wherein the alloy formed comprises Ni.
 12. The method of claim10, wherein the alloy formed comprises Co.
 13. The method of claim 10,wherein the alloy formed comprises Fe.
 14. The method of claim 4,wherein the alloy formed comprises a metalloid selected from the groupconsisting of B, Si, C and combinations thereof.
 15. The method of claim4, wherein the alloy formed further comprises a non-metal selected fromthe group consisting of Ge, P, S and combinations thereof.
 16. Themethod of claim 14, wherein the metalloid is C.
 17. The method of claim10, wherein the alloy formed comprises an element selected from thegroup consisting of C, Si, and combinations thereof.
 18. The method ofclaim 15, wherein the alloy formed further comprises Ge.
 19. The methodof claim 5, wherein the alloy is contacted with a flux material selectedfrom the group consisting of B₂O₃, glass, calcium oxide, barium oxide,aluminum oxide, magnesium oxide, lithium oxide, and combinationsthereof.
 20. The method of claim 19, wherein the alloy is contacted witha flux material selected from the group consisting B₂O₃, glass, andcombinations thereof. 21.-41. (canceled)